Method for manufacturing a hot dip galvanized and galvannealed steel sheet having excellent elongation properties

ABSTRACT

There is provided a hot-dip galvanized steel sheet and a hot-dip galvannealed steel sheet, which have excellent elongation properties, and methods for manufacturing the hot-dip galvanized steel sheet and the hot-dip galvannealed steel sheet. The present disclosure relates to a hot-dip galvanized steel sheet in which a hot-dip galvanized layer is formed on a surface of abase steel sheet, the hot-dip galvanized steel sheet having excellent elongation properties and being characterized by the composition and the microstructure thereof.

CROSS-REFERENCE TO RELATED APPLICATION

This a divisional patent application of U.S. patent application Ser. No.14/757,453, filed Dec. 23, 2015 which in turn claims the benefit ofKorean Patent Application No. 10-2014-0187935 filed on Dec. 24, 2014,the disclosures of which are incorporated herein by reference.

BACKGROUND

The present disclosure relates to a high-strength steel sheet, and moreparticularly, to a hot-dip galvanized steel sheet and a hot-dipgalvannealed steel sheet, applicable to automotive body panels, and thelike, due to possessing excellent formability, and methods formanufacturing the hot-dip galvanized steel sheet and the hot-dipgalvannealed steel sheet.

With a focus on crashworthiness regulations and fuel efficiency invehicles, high-strength steels are being actively used in automotiveapplications in order to achieve requirements for high strength,together with reductions in the weight of automotive bodies, and, inline with these trends, high-strength steels are being applied toautomotive body panels. Currently, a 340 MPa grade bake hardening steelis generally used for automotive body panels. However, 490 MPa gradebake hardening steel is also being used in some applications, while theapplication of 590 MPa grade bake hardening steel is further expected inthe future.

While the use of such bake hardening steel sheets for body panels allowsfor weight reductions, and improves impact resistance, the use of suchsteel sheets may be disadvantageous, in that formability decreases asstrength increases. Thus, recently, customers have demanded steel sheetshaving a low yield ratio (YR=YS/TS) and excellent ductility in order tocompensate for insufficient workability in high-strength steel sheetsapplied to automotive body panels. Furthermore, above all, steel sheetsused for automotive body panels should have excellent surface qualities.However, it may be difficult to ensure desired surface qualities inplated steel sheets due to hardenability-imparting and easily oxidizableelements (e.g., Si, Mn etc.), which are added to obtain high strength.

On the other hand, excellent corrosion resistance is required for theproper use of such steel sheets in automotive applications, and hot-dipgalvanized steel sheets having excellent corrosion resistance have thusbeen conventionally used as steel sheets for automotive applications.These steel sheets are manufactured through continuous hot-dipgalvanizing equipment in which recrystallization annealing and platingare performed on the same line, and thus have an advantage in thathighly corrosion resistant steel sheets may be manufactured at low cost.Furthermore, hot-dip galvannealed steel sheets, which are obtainedthrough hot-dip galvanizing and reheating thereafter, are being widelyimplemented due to excellent weldability or formability in addition toexcellent corrosion resistance.

Therefore, in order to reduce the weight of and increase the workabilityof automotive body panels, the development of a high-strengthcold-rolled steel sheet having excellent formability is required, and inaddition thereto, the development of a high-strength hot-dip galvanizedsteel sheet having excellent corrosion resistance, weldability, andformability is also required.

As prior art in which workability of such a high-strength steel sheet isimproved, patent document 1 discloses a multi-phase steel sheetincluding martensite as a main component, and suggests a method formanufacturing a high-strength steel sheet in which fine Cu precipitateshaving particle diameters of 1-100 nm are dispersed in the structure toimprove workability. However, according to the disclosure of patentdocument 1, it is necessary to add an excess amount of Cu, e.g., 2-5% ofCu, in order to precipitate fine Cu particles so that red shortnesscaused by Cu may occur, and manufacturing costs may be excessivelyincreased.

Patent document 2 discloses a multi-phase steel sheet including ferriteas a primary phase, retained austenite as a secondary phase, and bainiteand martensite, which are low-temperature transformed phases, and amethod for improving the ductility and stretch-flangeability of thesteel sheet. However, according to the disclosure of patent document 2,large amounts of Si and Al are added to obtain the retained austenitephase, and therefore it is difficult to ensure plating quality andsurface qualities in steelmaking and continuous casting. In addition,there is a disadvantage in that the initial YS value is high due totransformation induced plasticity, and thus the yield ratio is high.

Patent document 3, which relates to a feature for providing ahigh-strength hot-dip galvanized steel sheet exhibiting goodworkability, discloses a steel sheet having a microstructure including acomposite of relatively soft ferrite and relatively hard martensite, anda manufacturing method for improving the elongation and Lankford value(R-value) of the steel sheet. According to this feature, however, alarge amount of Si is added, and thus it is difficult to obtainexcellent plating quality, and manufacturing costs are increased due tothe addition of large amounts of Ti and Mo.

RELATED ART DOCUMENTS

(Patent Document 1) JP Patent Application Laid-open Publication No.2005-264176

(Patent Document 2) JP Patent Application Laid-open Publication No.2004-292891

(Patent Document 3) KR Patent Application Laid-open Publication No.2002-0073564

SUMMARY

An aspect of the present disclosure may provide a hot-dip galvanizedsteel sheet and a hot-dip galvannealed steel sheet, which are suitablefor automotive body panels, and in which a design of an alloy andmanufacturing conditions are optimized to achieve a tensile strength of450-650 MPa and to exhibit excellent low yield ratio (YS/TS) properties,and particularly, to allow for significantly improved ductility againstthe yield ratio.

An aspect of the present disclosure may also provide methods formanufacturing the hot-dip galvanized steel sheet and the hot-dipgalvannealed steel sheet.

However, the embodiments of the present disclosure are not limitedthereto, and other technical embodiments not mentioned herein would beclearly understood by a person skilled in the art from the descriptionbelow.

According to an aspect of the present disclosure, there is provided ahot-dip galvanized steel sheet in which a hot-dip galvanized layer isformed on a surface of a base steel sheet, the hot-dip galvanized steelsheet having excellent elongation properties, wherein

the base steel sheet:

contains 0.02-0.08% of carbon (C), 1.3-2.1% of manganese (Mn), 0.3% orless of silicon (Si) (excluding 0%), 1.0% or less of chromium (Cr)(excluding 0%), 0.1% or less of phosphorous (P) (excluding 0%), 0.01% orless of sulfur (S) (excluding 0%), 0.01% or less of nitrogen (N)(excluding 0%), 0.02-0.06% of aluminum (sol. Al), 0.2% or less ofmolybdenum (Mo) (excluding 0%), 0.003% or less of boron (B) (excluding0%), and a remainder of Fe and other inevitable impurities, in wt %;

has a microstructure including 90% or more of ferrite, and a remainderof martensite and 3% or less of bainite, as defined by Equation 1;

contains a percentage of martensite (M %) having an average particlediameter of 5 μm or less occupying space in ferrite grain boundaries(including grain boundary triple points) of 90% or higher, as defined byEquation 2; and

at ¼t based on the thickness (t) of the sheet, the ratio of a Cconcentration (a) in the ferrite phase to a C concentration (b) in themartensite phase, i.e., (a)/(b), is 0.7 or less, and the ratio of an Mnconcentration (c) in the ferrite phase to an Mn concentration (d) in themartensite phase, i.e., (c)/(d), is 0.8 or less.B(%)={BA/(MA+BA)}×100  [Equation 1]

where BA is the area occupied by bainite, and MA is the area occupied bymartensite.M(%)={M _(gb)/(M _(gb) +M _(in))}×100  [Equation 2]

where M_(gb) is the amount of martensite at ferrite grain boundaries,and M_(in) is the amount of martensite within ferrite grains, themartensite having an average particle diameter of 5 μm or less.

In the base steel sheet, the ferrite phase may have an average grainsize of 4 μm or greater, and the area occupied by the ferrite phasehaving an average grain size of 7 μm or greater in the entire ferritephase may be 10% or higher.

According to another aspect of the present disclosure, there is provideda hot-dip galvannealed steel sheet in which a hot-dip galvannealed layeris formed on a surface of a base steel sheet, the hot-dip galvannealedsteel sheet having excellent elongation properties, wherein

the base steel sheet:

contains 0.02-0.08% of carbon (C), 1.3-2.1% of manganese (Mn), 0.3% orless of silicon (Si) (excluding 0%), 1.0% or less of chromium (Cr)(excluding 0%), 0.1% or less of phosphorous (P) (excluding 0%), 0.01% orless of sulfur (S) (excluding 0%), 0.01% or less of nitrogen (N)(excluding 0%), 0.02-0.06% of aluminum (sol. Al), 0.2% or less ofmolybdenum (Mo) (excluding 0%), 0.003% or less of boron (B) (excluding0%), and a remainder of Fe and other inevitable impurities, in wt %;

has a microstructure including 90% or more of ferrite, and a remainderof martensite and 3% or less of bainite, as defined by Equation 1;

contains a percentage of martensite (M %) having an average particlediameter of 5 μm or less occupying space in ferrite grain boundaries(including grain boundary triple points) of 90% or higher, as defined byEquation 2; and

at ¼t based on the thickness (t) of the sheet, the ratio of a Cconcentration (a) in the ferrite phase to a C concentration (b) in themartensite phase, i.e., (a)/(b), is 0.7 or less, and the ratio of an Mnconcentration (c) in the ferrite phase to an Mn concentration (d) in themartensite phase, i.e., (c)/(d), is 0.8 or less.

In the base steel sheet, the ferrite phase may have an average grainsize of 4 μm or greater, and the area occupied by the ferrite phasehaving an average grain size of 7 μm or greater in the entire ferritephase may be 10% or higher.

According to another aspect of the present disclosure, there is provideda method for manufacturing a hot-dip galvanized steel sheet havingexcellent elongation properties, the method including:

preparing a steel slab having the above-described compositionalcomponents, and thereafter reheating the steel slab;

performing finish hot-rolling on the reheated steel slab in atemperature range of Ar3+50° C.-950° C., and thereafter coiling thefinish hot-rolled steel sheet at 450-700° C.;

performing cold-rolling on the coiled hot-rolled steel sheet with areduction ratio of 40-80%, and thereafter performing continuousannealing on the cold-rolled steel sheet in a temperature range of760-850° C.;

performing a first cooling on the continuous annealed steel sheet to atemperature range of 630-670° C. at an average cooling rate of 2-8°C./s, and thereafter performing a second cooling on the first cooledsteel sheet to a temperature range of Ms+20° C. to Ms+50° C. at anaverage cooling rate of 3-10° C./s; and

performing hot-dip galvanizing on the second cooled steel sheet in atemperature range of 440-480° C., and thereafter cooling the hot-dipgalvanized steel sheet to a temperature of Ms−100° C. or lower at anaverage cooling rate of 4° C./s or higher.

In the base steel sheet of the hot-dip galvanized steel sheet,

the microstructure may include 90% or more of ferrite, and a remainderof martensite and 3% or less of bainite, as defined by Equation 1;

the percentage of martensite (M %) having an average particle diameterof 5 μm or less occupying space in ferrite grain boundaries (includinggrain boundary triple points) may be 90% or higher, as defined byEquation 2; and

at ¼t based on the thickness (t) of the sheet, the ratio of a Cconcentration (a) in the ferrite phase to a C concentration (b) in themartensite phase, i.e., (a)/(b), may be 0.7 or less, and the ratio of anMn concentration (c) in the ferrite phase to an Mn concentration (d) inthe martensite phase, i.e., (c)/(d), may be 0.8 or less.

In the base steel sheet of the hot-dip galvanized steel sheet, theferrite phase may have an average grain size of 4 or greater, and thearea occupied by the ferrite phase having an average grain size of 7 μmor greater in the entire ferrite phase may be 10% or higher.

According to another aspect of the present disclosure, there is provideda method for manufacturing a hot-dip galvannealed steel sheet havingexcellent elongation properties, the method including:

preparing a steel slab having the above-described compositionalcomponents, and thereafter reheating the steel slab;

performing finish hot-rolling on the reheated steel slab in atemperature range of Ar3+50° C.-950° C., and thereafter coiling thefinish hot-rolled steel sheet at 450-700° C.;

performing cold-rolling on the coiled hot-rolled steel sheet with areduction ratio of 40-80%, and thereafter performing continuousannealing in which the cold-rolled steel sheet is subjected to a secondheating to a temperature range of 760−850° C. at an average heating rateof 2° C./s or less;

performing a first cooling on the continuous annealed steel sheet to atemperature range of 630-670° C. at an average cooling rate of 2-8°C./s, and thereafter performing a second cooling on the first cooledsteel sheet to a temperature range of Ms+20° C. to Ms+50° C. at anaverage cooling rate of 3-10° C./s; and

performing hot-dip galvanizing on the second cooled steel sheet in atemperature range of 440-480° C., and thereafter performing alloyingheat treatment on the hot-dip galvanized steel sheet, and then coolingthe alloying heat-treated steel sheet to a temperature of Ms−100° C. orlower at an average cooling rate of 4° C./s or higher.

In the base steel sheet of the hot-dip galvannealed steel sheet,

the microstructure may include 90% or more of ferrite, and a remainderof martensite and 3% or less of bainite, as defined by Equation 1;

the percentage of martensite (M %) having an average particle diameterof 5 μm or less occupying space in ferrite grain boundaries (includinggrain boundary triple points) may be 90% or higher, as defined byEquation 2; and

at ¼t based on the thickness (t) of the sheet, the ratio of a Cconcentration (a) in the ferrite phase to a C concentration (b) in themartensite phase, i.e., (a)/(b), may be 0.7 or less, and the ratio of anMn concentration (c) in the ferrite phase to an Mn concentration (d) inthe martensite phase, i.e., (c)/(d), may be 0.8 or less.

In the base steel sheet of the hot-dip galvannealed steel sheet, theferrite phase may have an average grain size of 4 μm or greater, and thearea occupied by the ferrite phase having an average grain size of 7 μmor greater in the entire ferrite phase may be 10% or higher.

BRIEF DESCRIPTION OF DRAWINGS

The above and other aspects, features and other advantages of thepresent disclosure will be more clearly understood from the followingdetailed description taken in conjunction with the accompanyingdrawings, in which:

FIG. 1 is a graph showing concentration ratios of C and Mn in a ferritephase and a martensite phase at ¼t based on the thickness (t) of abasesteel sheet of a hot-dip galvanized steel sheet manufactured accordingto an embodiment of the present disclosure; and

FIG. 2 is an image showing a result of microstructural observation at ¼tbased on the thickness (t) of the base steel sheet in FIG. 1.

DETAILED DESCRIPTION

Exemplary embodiments of the present disclosure will now be described indetail with reference to the accompanying drawings.

The disclosure may, however, be exemplified in many different forms andshould not be construed as being limited to the specific embodiments setforth herein. Rather, these embodiments are provided so that thisdisclosure will be thorough and complete, and will fully convey thescope of the disclosure to those skilled in the art.

As a result of intensive investigation for providing a steel sheet whichhas both strength and ductility, and thus exhibits excellent formabilityto be suitable for automotive body panels, the present inventors foundthat a multi-phase steel sheet satisfying desired properties can beprovided by optimizing manufacturing conditions together with alloydesign, and finally achieved the present disclosure.

First, the hot-dip galvanized steel sheet and the hot-dip galvannealedsteel sheet, which have excellent elongation properties, according tothe present disclosure will be described in detail.

The hot-dip galvanized steel sheet according to the present disclosurecontains 0.02-0.08% of carbon (C), 1.3-2.1% of manganese (Mn), 0.3% orless of silicon (Si) (excluding 0%), 1.0% or less of chromium (Cr)(excluding 0%), 0.1% or less of phosphorous (P) (excluding 0%), 0.01% orless of sulfur (S) (excluding 0%), 0.01% or less of nitrogen (N)(excluding 0%), 0.02-0.06% of acid-soluble aluminum (sol. Al), 0.2% orless of molybdenum (Mo) (excluding 0%), 0.003% or less of boron (B)(excluding 0%), and a remainder of Fe and other inevitable impurities,in wt %.

Hereinafter, reasons for limiting the alloying composition of thehot-dip galvanizing steel sheet according to the present disclosure asabove will be described in detail. The content of each component refersto wt % unless otherwise specifically stated.

C: 0.02-0.08%

Carbon (C) is an important component in manufacturing a multi-phasesteel sheet, and is a beneficial element in terms of strength by formingmartensite, which is one of the dual-phase structures. In general, thehigher C content facilitates the formation of martensite, and is thusadvantageous in manufacturing a multi-phase steel. However, it isrequired that the C content is controlled at an appropriate level so asto control desired strength and yield ratio (YS/TS).

In particular, the higher C content tends to increase the yield ratio ofsteel because bainite transformation also occurs during cooling afterannealing. In the present disclosure, it is important, if possible, tominimize bainite formation and form martensite at an appropriate level,thereby obtaining desired material properties.

Therefore, it is preferable that the C content is controlled to be 0.02%or more. When the C content is less than 0.02%, it is difficult toobtain 450 MPa grade strength aimed in the present disclosure and toform martensite at an appropriate level. Conversely, when the C contentis more than 0.08%, the formation of inter-granular bainite isaccelerated during cooling after annealing, thereby increasing yieldstrength and the yield ratio (YS/TS), so that bending and surfacedefects are prone to occur in the formation of automotive components.Therefore, in the present disclosure, it is preferable that the Ccontent is controlled to be 0.02-0.08%, and more preferably, the Ccontent is controlled to be 0.03-0.06% to obtain appropriate strength.

Mn: 1.3-2.1%

Manganese (Mn) is a hardenability improving element in a multi-phasesteel sheet, and is an important element particularly in formingmartensite. In conventional solid-solution strengthened steel, Mn iseffective in improving strength due to the solid-solution strengtheningeffect, and precipitate S, which is inevitably added to steel, into MnS,thus playing an important role to suppress plate rupture and hightemperature embrittlement, which are caused by S during hot rolling.

In the present disclosure, it is preferable that 1.3% or more of Mn isadded. When the Mn content is less than 1.3%, martensite cannot beformed, so that it is difficult to manufacture a multi-phase steel.Conversely, when the Mn content is more than 2.1%, an excess amount ofmartensite is formed, thus causing material instability, and Mn-Band (Mnoxide band) is formed in the structure, so that the risk of work cracksand plate rupture is increased. Furthermore, Mn oxides are released tothe surface during annealing, thereby significantly deterioratingcoatability. Therefore, in the present disclosure, it is preferable thatthe Mn content is limited to 1.3-2.1%, and more preferably, the Mncontent is limited to 1.4-1.8%.

Cr: 1.0% or less (excluding 0%)

Chromium (Cr) is a component having similar properties to theabove-described Mn, and is added to improve hardenability and to obtainhigh strength in steel. Cr is effective in forming martensite, and formscoarse Cr carbides such as Cr₂₃C₆ during a hot rolling process, therebyprecipitating C dissolved in steel below an appropriate level, and thussuppressing yield point elongation (YP-EI). Therefore, Cr is abeneficial element in manufacturing a multi-phase steel having a lowyield ratio. Furthermore, Cr minimizes a decrease in elongation againstan increase in strength, and is also beneficial in manufacturing amulti-phase steel having high ductility.

In the present disclosure, the aforementioned Cr facilitates theformation of martensite through improving hardenability. However, whenthe Cr content is more than 1.0%, the fraction of forming martensiteexcessively increases, thereby causing decreases in strength andelongation.

Therefore, in the present disclosure, it is preferable that the Crcontent is limited to 1.0% or less, excluding 0% in consideration ofbeing inevitably added during manufacturing.

Si: 0.3% or less (excluding 0%)

In general, silicon (Si) is an element that forms retained austenite atan appropriate level during cooling after annealing, and thussignificantly contributing to improving elongation. However, this iseffective when the C content is high (e.g., about 0.6%). Furthermore, itis known that the aforementioned Si acts to improve the strength ofsteel through the solid-solution strengthening effect, or, above anappropriate level, improves surface properties of hot-dip galvanizedsteel sheets.

In the present disclosure, the Si content is limited to 0.3% or less(excluding 0%) for obtaining sufficient strength and improvingelongation. Even without the addition of Si, there is no significantproblem in terms of properties. However, 0% is excluded in considerationof being inevitably added during manufacturing. When the Si content ismore than 0.3%, surface properties of plated steel sheets aredeteriorated, and there is almost no effect in forming a multi-phasesteel.

P: 0.1% or less (excluding 0%)

Phosphorous (P) in steel is the most beneficial element in obtainingsufficient strength without significantly deteriorating formability.However, it is problematic in that when P is excessively added, thepossibility of brittle fracture significantly increases, therebyincreasing the possibility of plate rupture during slab hot rolling, andP acts as an element that deteriorates surface properties of platedsteel sheets.

Therefore, in the present disclosure, the P content is limited to 0.1%,excluding 0% in consideration of being inevitably added.

S: 0.01% or less (excluding 0%)

Sulfur (S) is an impurity element inevitably added in steel, and it isimportant to control the S content to be as low as possible.Particularly, since S in steel increases the possibility of redshortness, it is preferable that the S content is controlled to be 0.01%or less. However, 0% is excluded in consideration of being inevitablyadded during manufacturing.

N: 0.01% or less (excluding 0%)

Nitrogen (N) is an impurity element inevitably added in steel. Althoughit is important to control the N content to be as low as possible, steelrefining costs dramatically increase to this end. Therefore, it ispreferable that the N content is controlled to be 0.01% or less in whichoperating conditions are feasible. However, 0% is excluded inconsideration of being inevitably added during manufacturing.

sol. Al: 0.02-0.06%

Acid-soluble aluminum (sol. Al) is an element that is added for grainrefinement and deoxidation in steel. When the sol. Al content is lessthan 0.02%, Al-killed steel cannot be produced in a conventional stablestate. Conversely, when the sol. Al content is more than 0.06%, it isadvantageous in terms of strength improvement, but excess amounts ofinclusions are formed during continuous casting in steelmaking, therebyincreasing the possibility of surface defects in hot-dip galvanizedsteel sheets and causing an increase in manufacturing costs. Therefore,in the present disclosure, it is preferable that the sol. Al content iscontrolled to be 0.02-0.06%.

Mo: 0.2% or less (excluding 0%)

Molybdenum (Mo) is an element that is added to delay transformation ofaustenite into pearlite and to accomplish refinement of ferrite andstrength improvement. Mo improves hardenability of steel and allows forforming fine martensite in grain boundaries, and is thus capable ofcontrolling the yield ratio. However, as an expensive element, thehigher content is unfavorable in manufacturing. Therefore, it ispreferable that the Mo content is appropriately controlled.

In order to obtain the above-described effects, it is preferable that Mnis added up to 0.2%. A Mn content of more than 0.2% leads to a sharpincrease in alloy costs, thereby reducing economical efficiency ratherthan reducing ductility of steel. Although, in the present disclosure,the optimum content of Mo is 0.05%, there is no difficulty in obtainingdesired properties even without essential addition. However, 0% isexcluded in consideration of being inevitably added duringmanufacturing.

B: 0.003% or less (excluding 0%)

Boron (B) in steel is an element that is added to prevent secondaryworking brittleness caused by the addition of P. A B content of morethan 0.003% leads to a decrease in elongation. Therefore, the B contentis controlled to be 0.003% or less, excluding 0% in consideration ofbeing inevitably added.

The steel sheet according to the present disclosure may contain aremainder of Fe and other inevitable impurities, in addition to theaforementioned components.

In the hot-dip galvanized steel sheet and the hot-dip galvannealed steelsheet, which satisfy the above-described compositions, according to thepresent disclosure, it is preferable that the base steel sheet has amicrostructure including ferrite as a primary phase, and a remainder ofmartensite. In this case, some of the microstructure may includebainite, and it is preferable that the amount of bainite is, ifpossible, minimized, or that bainite is absent.

Therefore, it is preferable that the base steel sheet of the hot-dipgalvanized steel sheet according to the present disclosure has amicrostructure including, in area %, 90% or more of ferrite, and aremainder of martensite and 3% or less of bainite (B), as defined byEquation 1.

It is preferable that, at ¼t based on the total thickness (t) of thebase steel sheet, a ferrite fraction is 90% or more, and the fraction ofthe dual-phase structure including the remainder of martensite andbainite is 1-10%. When the fraction of the dual-phase structure is lessthan 1%, it is difficult to form a multi-phase steel, and thus difficultto obtain a steel sheet having a low yield ratio. Conversely, when thefraction of the dual-phase structure is more than 10%, the strengthbecomes too high to obtain desired workability. According toexperimental results of the present inventors, a more preferablefraction of a martensite structure is 2-5% at ¼t based on the thicknessof the base steel sheet. This is the optimum condition for obtaining anexcellent low yield ratio and sufficient ductility by controlling theoptimum content of fine martensite. Furthermore, as shown in Equation 1below, although bainite may be absent, 3% or less of bainite ispreferable when bainite is inevitably formed. When the bainite contentis more than 3%, C concentration around bainite increases, thusdeteriorating ductility and increasing the yield ratio, thereby beinginappropriate for the present disclosure.B (%)={BA/(MA+BA)}×100  [Equation 1]

where BA is the area occupied by bainite, and MA is the area occupied bymartensite.

In the present disclosure, it is important that the area ratio ofbainite in the entire dual-phase structure is controlled to be low. Thisis because, in the case of bainite compared to martensite, solute atomsin bainite grains, i.e., C and N, are easily trapped in dislocations,resulting in an impediment of dislocation movement and discontinuousyielding behavior, thereby significantly increasing the yield ratio.

Therefore, when the area ratio of bainite in the entire dual-phasestructure is 3% or less, the yield ratio prior to skin pass rolling maybe controlled to be 0.57 or less, and the yield ratio may be controlledat an appropriate level by subsequent skin pass rolling. When the arearatio of bainite is higher than 3%, the yield ratio prior to skin passrolling is higher than 0.57, so that it is difficult to manufacture amulti-phase steel sheet having a low yield ratio, and thus causing adecrease in ductility.

Furthermore, in the hot-dip galvanized steel sheet and the hot-dipgalvannealed steel sheet according to the present disclosure, it ispreferable that the ratio (M %) occupied by martensite having an averageparticle diameter of 5 μm or less in ferrite grain boundaries (includinggrain boundary triple points), which is defined by Equation 2 below, is90% or higher. That is, when the fine martensite having an averageparticle diameter of 5 μm or less exists mainly in ferrite grainboundaries compared to within ferrite grains, it is advantageous inimproving ductility while a low yield ratio is being maintained.M (%)={M _(gb)/(M _(gb) +M _(in))}×100  [Equation 2]

where M_(gb) is the amount of martensite in ferrite grain boundaries,and M_(in) is the amount of martensite within ferrite grains, themartensite having an average particle diameter of 5 μm or less.

Thus, when the ratio occupied by martensite in ferrite grain boundariesis 90% or higher, the yield ratio prior to skin pass rolling may becontrolled to be 0.55 or less, and the yield ratio may be controlled atan appropriate level by subsequent skin pass rolling. When the ratiooccupied by martensite is less than 90%, martensite formed within grainsincreases yield strength during tensile strain, so that the yield ratioincreases and cannot be controlled through skin pass rolling. Inaddition, elongation decreases because martensite within grainssignificantly impedes dislocation movement during formation, therebyweakening ductility and thus causing a decrease in elongation. Also,elongation decreases because a large amount of martensite is formedwithin ferrite grains, thereby generating too many dislocations, andthus impeding the movement of mobile dislocations during forming. Thepresent disclosure is devised so as to maximize F content and thusachieve ductility.

Furthermore, in the hot-dip galvanized steel sheet and the hot-dipgalvannealed steel sheet according to the present disclosure, it ispreferable that at ¼t based on the thickness of the base steel sheet,the ratio of a C concentration (a) in the ferrite phase to a Cconcentration (b) in the martensite phase, i.e., (a)/(b), is 0.7 orless, and the ratio of an Mn concentration (c) in the ferrite phase toan Mn concentration (d) in the martensite phase, i.e., (c)/(d), is 0.8or less.

These features are very important in the present disclosure, and thetechnical significance thereof is that a multi-phase steel havingexcellent ductility is provided even in the same ferrite content bycontrolling concentration ratios of C and Mn in ferrite and martensitein a matrix structure, such that concentration ratios of C and Mn inferrite are controlled to be as low as possible as compared tomartensite. At ¼t based on the thickness of the base steel sheet, whenthe ratio of a C concentration (a) in the ferrite phase to a Cconcentration (b) in the martensite phase, i.e., (a)/(b), is 0.7 or lessand the Mn concentration ratio is less than 0.8, the softening propertyof ferrite is improved, and better ductility may be obtained. However,when the C concentration ratio is higher than 0.7 or the Mnconcentration ratio is higher than 0.8, there is almost no difference inC and Mn concentration ratios between martensite and ferrite, so that itis not possible to obtain required ductility. This may be developed bycomponent designs and characterization of operating conditions.

Meanwhile, in the hot-dip galvanized steel sheet and the hot-dipgalvannealed steel sheet according to the present disclosure, it ispreferable that, in the base steel sheet, the ferrite phase has anaverage grain size of 4 μm or greater, and the area occupied by theferrite phase having an average grain size of 7 μm or greater in theentire ferrite phase is controlled to be 10% or higher. When the ferritephase has the larger grain size and more uniform grains, it is possibleto manufacture a steel sheet having excellent ductility. When theaverage grain size is less than 4 μm, it is not possible to obtain therequired ductility. Also, when the area occupied by the ferrite phasehaving an average grain size of 7 μm or greater in the entire ferritephase is less than 10%, ferrite grain coarsening cannot be performed,thereby being problematic in obtaining sufficient ductility.

Next, the method for manufacturing a hot-dip galvanized steel sheet anda hot-dip galvannealed steel sheet having excellent elongationproperties according to the present disclosure will be described indetail.

First, in the present disclosure, a steel slab having compositionalcomponents as described above is prepared and then reheated. Such areheating step is performed in order to perform a subsequent hot-rollingstep without a hitch, and to fully obtain the desired properties of asteel sheet. The present disclosure is not particularly limited to suchreheating conditions, but general conditions may be employed. As anexample, the reheating step may be performed in in a temperature rangeof 1100-1300° C.

Subsequently, in the present disclosure, the reheated steel slab issubjected to finish hot-rolling in a temperature range of Ar3+50°C.-950° C. In the present disclosure, it is preferable that the reheatedsteel slab is subjected to finish hot-rolling in a temperature range ofAr3+50° C.-950° C. which is defined by Equation 3 below. Fundamentally,it is advantageous to perform the finish hot-rolling in an austeniticsingle phase region. This is because more uniform deformation is appliedto a structure basically composed of single phase grains, andhomogeneity in the structure may thus be increased. When the finishhot-rolling temperature is below Ar3+50° C., rolling is highly likely tobe performed in a ferrite+austenite dual-phase region, thereby causingpoor material homogeneity. Conversely, when the finish hot-rollingtemperature is above 950° C., during cooling after hot-rolling, coilwarping may occur due to material inhomogeneity caused by the formationof abnormal coarse grains.Ar3=910−310*C−80*Mn−20*Cu−15*Cr−55*Ni−80*Mo  [Equation 3]

where Ar3 is a theoretical temperature.

Then, in the present disclosure, the finish hot-rolled sheet is coiledat 450-700° C. When the coiling temperature is lower than 450° C., anexcess amount of martensite or bainite is formed, thus causing anexcessive increase in strength, so that problems such as shape defectsmay be caused by a load during subsequent cold-rolling. Conversely, whenthe coiling temperature is above 700° C., there is a problem of severesurface enrichment of elements such Si, Mn, and B in steel, which reducehot-dip galvanizing wettability. Therefore, it is preferable that thecoiling temperature is controlled to be 450-700° C. in consideration ofthese problems. Then, the coiled hot-rolled sheet may be subjected topickling under general conditions.

Subsequently, in the present disclosure, the coiled hot-rolled steelsheet is subjected to cold-rolling with a reduction ratio of 40-80%. Itis preferable that the cold-rolling is performed with a reduction ratioof 40-80%. When the cold-rolling reduction ratio is less than 40%, it isdifficult to obtain a desired thickness and straighten the steel sheet.Conversely, when the cold-rolling reduction ratio is greater than 80%,it is highly likely that cracks are generated at the edge of the steelsheet, which may result in an overload on cold-rolling.

Subsequently, the cold-rolled steel sheet prepared according to theabove steps is preferably subjected to continuous annealing in atemperature range of 760-850° C. The continuous annealing is performedin a continuous galvannealing furnace.

The purpose of the continuous annealing step is to performrecrystallization, form ferrite and austenite recrystallization, anddistribute carbon. When the continuous annealing temperature is below760° C., full recrystallization is not achieved, and it is difficult tosufficiently form austenite, so that it becomes difficult to obtain thestrength intended in the present disclosure. Conversely, when thecontinuous annealing temperature is above 850° C., it is problematic inthat productivity decreases, and austenite is excessively formed, thusreducing ductility due to bainite after cooling. Therefore, it ispreferable that the continuous annealing temperature range is controlledto be 760-850° C. in consideration of these problems. More preferably,the continuous annealing is performed in a temperature range of 770-810°C.

Although these temperature ranges correspond to dual-phase regions(ferrite+austenite), it is preferable that the continuous annealing isperformed at a temperature in which ferrite regions are included as muchas possible. The higher amount of initial ferrite at an annealingtemperature in dual-phase regions allows for facilitating grain growthafter annealing, thus improving ductility. Lowering a martensite starttemperature (Ms) with an increase in the degree of C enrichment inaustenite enables the formation of martensite during final cooling aftersubsequent hot-dip galvanizing in a pot, and fine and uniform martensiteis thus largely distributed in grains, thereby being capable of themanufacturing of a steel sheet having excellent ductility and low yieldratio. The lower the secondary heating temperature range, the morefavorably C in ferrite diffuses into austenite (higher degree of Csaturation in austenite than ferrite), and the higher content ofaustenite having a higher degree of C enrichment facilitates theformation of fine martensite, thereby being capable of the manufacturingof a steel sheet having high ductility.

Furthermore, in the present disclosure, the continuous annealed steelsheet is subjected to a first cooling at an average cooling rate of 2-8°C./s. In the present disclosure, the higher first cooling temperatureand the lower first cooling rate allow for more uniform ferrite and ahigher coarsening tendency, thus being advantageous in terms ofductility. Furthermore, during the first cooling, enough time is givenfor C, albeit a small amount, to diffuse into austenite. This has animportant significance in the present disclosure. In more detail, in adual-phase region, C always dynamically diffuses into austenite,generally, having high degree of C enrichment, and diffusion velocityincreases with increases in temperature and time. Therefore, the firstcooling temperature is important. When the first cooling temperature isbelow 630° C., which is too low of a temperature, diffusion activity ofC is too low to sufficiently diffuse into austenite, and C concentrationin ferrite thus increases. Therefore, in a final material, a Cconcentration gradient between ferrite and martensite phases becomes 0.7or higher, or a Mn concentration gradient therebetween becomes 0.8 orhigher, thus being disadvantageous in terms of ductility. That is, thepresent disclosure is characterized in that C and Mn concentrationratios in ferrite are controlled to be as low as possible in order tofacilitate ferrite softening and thus manufacture steel having excellentductility. Conversely, when the first cooling temperature is above 670°C., it is advantageous in terms of the aforementioned characteristics,but is problematic in that an overly rapid cooling rate may be requiredduring cooling.

Also, it is preferable that the first cooling rate is limited to 2-8°C./s. This is because when the first cooling rate is less than 2° C./s,it is problematic in terms of productivity due to an overly low coolingrate, and when the first cooling rate is higher than 8° C./s, there isinsufficient time for C to diffuse into austenite.

Subsequently, in the present disclosure, the first cooled steel sheet issubjected to a second cooling to a temperature range of Ms+20° C.-Ms+50°C. at an average cooling rate of 3-10° C./s. Here, Ms may be defined byEquation 4 below.Ms (° C.)=539−423C−30.4Mn−12.1Cr−17.7Ni−7.5Mo  [Equation 4]

where Ms is a theoretical temperature for forming a M phase.

According to the present disclosure, when the martensite phase is formedbefore passing through 440-480° C., which is a temperature range ingeneral hot-dip galvanizing pots, finally, the martensite phase is proneto coarsening, and it is not possible to obtain a low yield ratio.Therefore, in the present disclosure, the second cooling temperaturerange is limited to Ms+20° C.-Ms+50° C., and it is required that thesecond cooling rate is as low as possible in this temperature conditionto suppress the formation of the martensite phase. When the secondcooling temperature is below Ms+20° C., a martensite phase may beformed. Conversely, when the second cooling temperature is above Ms+50°C., a subsequent cooling rate becomes rather high, and the martensitephase is highly likely to form before subsequent dipping into the pot.

Meanwhile, it is preferable that the second cooling rate is limited to3-10° C./s. This is because when the second cooling rate is less than 3°C./s, the martensite phase is not formed, but is problematic in terms ofproductivity, and when the second cooling rate is higher than 10° C./s,threading speed increases as a whole, and problems such as warping in aplate shape may thus be caused.

Subsequently, in the present disclosure, the second cooled steel sheetis subjected to a hot-dip galvanizing treatment in a temperature rangeof 440-480° C., and thereafter, is cooled to a temperature of Ms−100° C.or lower at an average cooling rate of 4° C./s or higher.

In the present disclosure, the hot-dip galvanizing treatment may beperformed by dipping the second cooled steel sheet into the pot in ageneral temperature range of 440-480° C. Although the present disclosureis not limited to these specific hot-dip galvanizing conditions, it ispreferable that the average cooling rate, with which the second cooledsteel sheet is cooled to a pot temperature in the above temperaturerange, is set to be 4-8° C./s. By controlling the average cooling rateto be 4-8° C./s, it is possible that a martensite structure in the steelsheet is not formed before the steel sheet reaches the pot.Specifically, when the cooling rate is lower than 4° C./s, martensite isnot formed, but is inappropriate due to poor productivity. Conversely,when the cooling rate is higher than 8° C./s, some martensite andbainite are formed within grains, thus increasing yield strength anddeteriorating ductility.

Subsequently, in the present disclosure, the hot-dip galvanized steelsheet is cooled to a temperature of Ms−100° C. or lower at an averagecooling rate of 4° C./s or higher, so that a hot-dip galvanized steelsheet having fine martensite may be manufactured in the final step. Intemperature conditions above Ms−100° C., fine martensite phases cannotbe obtained unless the cooling rate is very high, and shape defects inthe sheet may occur.

Therefore, in the present disclosure, the hot-dip galvanized steel sheetis cooled to a temperature of Ms−100° C. or lower at an average coolingrate of 4° C./s or higher. When the cooling rate is less than 4° C./s,due to an overly low cooling rate, martensite is irregularly formed ingrain boundaries or within grains, and the formation ratio ofinter-granular martensite to intra-granular martensite is too low, sothat steel having a low yield ratio cannot be manufactured andproductivity is also deteriorated.

Meanwhile, in the present disclosure, the above-described hot-dipgalvanized steel sheet may be subjected to a subsequent reheatingtreatment at a general heat treating temperature, and then finallycooled to a temperature of Ms−100° C. or lower at an average coolingrate of 4° C./s or higher to manufacture a hot-dip galvannealed steelsheet. Other conditions are the same as in the case of theabove-described hot-dip galvanized steel sheet.

Hereinafter, the present disclosure will be described in detail by wayof example.

Steel slabs having compositional components shown in Table 1 below wereprepared, and thereafter, hot-dip galvanized steel sheets weremanufactured using manufacturing processes as shown in Table 2 below. InTable 1 below, inventive steels 1-2 and 4-5 were used to manufacturehot-dip galvannealed steel sheets (GA), and inventive steels 3 and 6were used to manufacture hot-dip galvanized steel sheets (GA).

Properties of hot-dip galvanized steel sheets manufactured as above wereevaluated, and the results are shown in Table 3 below. The presentdisclosure aims to manufacture a steel sheet having a yield ratio of0.57 or less and a hole expansion ratio (HER) of 80% or higher in astate without skin pass rolling.

A tensile test for each specimen was performed in the C direction usingthe JIS standard, a matrix structure was analyzed at ¼t based thethickness of an annealed steel sheet, and microstructural fractions wereobtained using the results. Specifically, the area ratios of martensiteand bainite were first calculated using an optical microscope throughLepelar etching, and then the fractions thereof were accurately measuredthrough a Count Point operation after observation using a SEM(magnification: ×3,000)

Meanwhile, in order to obtain C and Mn concentration ratios in ferriteand martensite in the matrix structure of a base steel sheet, asputtering cut was performed from the surface of a plated layer to ¼tinside the base steel without structural damage using a Focus Ion Beam(FIB) system. Subsequently, the concentration ratios of C and Mn in eachphase were quantitatively evaluated by Line and Point scanning in EnergyDispersive Spectroscopy (EDS) analysis using TEM.

Steels having compositions shown in Table 1 below were prepared underthe conditions shown in Table 2 below, and thereafter, propertiesthereof were confirmed. As desired material properties in the presentdisclosure, the target is to obtain a yield ratio of 0.57 or less and ahole expansion ratio (HER) of 80% or higher in a state without skin passrolling.

A tensile test for each specimen was performed in the C direction usingthe JIS standard, and microstructural fractions were obtained at ¼tbased on the thickness of an annealed steel sheet using an opticalmicroscope. M and B phases were observed through Lepelar etching. Thearea ratios of martensite and bainite were first calculated using anoptical microscope through Lepelar etching, and then the fractionsthereof were accurately measured through a Count Point operation afterobservation using a SEM (magnification: ×3,000).

In order to obtain C and Mn concentration ratios in ferrite andmartensite in the structure, a sputtering cut was performed from thesurface of a plated layer to ¼t inside the base steel without structuraldamage using a Focus Ion Beam (FIB) system. A 10 mm hole was created inthe specimen by milling, the specimen was then pushed up from the bottomusing a cone-type punch until surface cracking was initiated, and thehole expansion ratio was obtained by calculating the initial holediameter ratio before the cracking.

In Table 3 below, the yield ratio (7) is a value measured before skinpass rolling.

TABLE 1 Composition (wt %) Item C Mn Si Cr P S N sol. Al Mo B Inventive0.023 1.75 0.05 0.81 0.06 0.006 0.003 0.03 0.15 0.0006 Steel 1 Inventive0.041 1.72 0.04 0.48 0.05 0.005 0.003 0.04 0.12 0.0009 Steel 2 Inventive0.053 1.55 0.11 0.42 0.03 0.007 0.004 0.05 0.13 0.0015 Steel 3 Inventive0.056 1.56 0.16 0.83 0.04 0.004 0.003 0.041 0.15 0.0021 Steel 4Inventive 0.068 1.48 0.21 0.87 0.02 0.003 0.005 0.052 0.18 0.0008 Steel5 Inventive 0.076 1.38 0.25 0.08 0.03 0.004 0.008 0.025 0.08 0.0012Steel 6 Comparative 0.095 1.22 0.6 1.16 0.13 0.006 0.003 0.04 0.45 0.004Steel 1 Comparative 0.091 1.26 0.8 1.21 0.12 0.007 0.005 0.05 0.380.0041 Steel 2

TABLE 2 Manufacturing Process Conditions Finish Cold-rolling FirstSecond Pot Final Reheating Rolling Coiling Reduction Annealing CoolingCooling Cooling Cooling Temperature Temperature Temperature RatioTemperature Rate Rate Rate Rate Item (SRT ° C.) (FT ° C.) (° C.) (%) (°C.) (° C./s) (° C./s) (° C./s) (° C./s) Remarks Inventive 1182 882 58046 768 2.5 4.2 4.5 4.5 Example 1 Steel 1186 895 556 52 769 2.8 4.5 4.6 .5.7 Example 2 1 Inventive 1189 912 465 63 775 3.5 3.8 5.1 6.2 Example 3Steel 1186 921 472 65 779 3.6 3.5 5.5 6.3 Example 4 2 Inventive 1201 891682 71 810 4.8 6.3 6.3 9.2 Example 5 Steel 1203 896 678 72 815 4.9 6.26.2 9.6 Example 6 3 Inventive 1195 935 580 78 751 5.6 8.1 7.8 5.3Comparative Steel Example 1 4 1196 942 585 79 821 5.8 8.6 7.5 7.8Example 7 Inventive 1186 928 630 69 855 6.8 9.4 9.2 10.2 ComparativeSteel Example 2 5 1189 912 632 65 839 6.2 12.1 10.8 9.2 ComparativeExample 3 Inventive 1210 897 682 38 841 7.5 8.5 9.2 5.2 ComparativeSteel Example 4 6 1206 888 675 68 836 9.8 7.8 9.5 8.9 ComparativeExample 5 Comparative 1203 897 660 72 802 2.8 6.5 11.5 5.3 ComparativeSteel Example 6 1 Comparative 1196 892 672 75 803 3.5 6.5 6.8 5.2Comparative Steel Example 7 2 1187 885 683 78 779 4.1 7.8 8.3 3.8Comparative Example 8

TABLE 3 {circle around (1)} {circle around (2)} {circle around (3)}{circle around (6)} {circle around (7)} M Fraction B Fraction M Ratio{circle around (4)} {circle around (5)} F Ratio Yield Elongation TS Item(%) (%) (%) (a/b) (c/d) (%) Ratio (%) (MPa) Remarks Inventive 2.5 1.5 920.65 0.71 12 0.57 37 475 Example 1 Steel 1.8 1.7 91.5 0.63 0.69 13 0.5638 468 Example 2 1 Inventive 3.5 0.7 90.6 0.58 0.65 12 0.57 36 509Example 3 Steel 3.4 0.5 92.1 0.59 0.59 11 0.56 35 508 Example 4 2Inventive 4.2 0.3 92.3 0.62 0.63 12 0.57 33 518 Example 5 Steel 5.1 0.191.5 0.63 0.75 14 0.57 33 513 Example 6 3 Inventive 1.8 3.5 88 0.75 0.8313 0.63 23 612 Comparative Steel Example 1 4 6.8 2.1 90.5 0.58 0.71 120.57 30 632 Example 7 Inventive 3.1 0.3 87 0.55 0.63 8.2 0.65 25 532Comparative Steel Example 2 5 10.2 1.2 78 0.75 0.83 6.8 0.56 28 535Comparative Example 3 Inventive 1.8 0.5 93 0.72 0.82 7.8 0.65 32 532Comparative Steel Example 4 6 9.8 1.5 76 0.71 0.83 8.1 0.56 26 532Comparative Example 5 Comparative 4.1 3.5 77 0.77 0.82 7.5 0.72 26 554Comparative Steel Example 6 1 Comparative 4.3 3.2 81 0.75 0.83 7.2 0.7125 555 Comparative Steel Example 7 2 4.2 3.3 83 0.74 0.83 7.9 0.68 26541 Comparative Example 8

In Table 3,

{circle around (1)} is the fraction of martensite (M) in a structure(%),

{circle around (2)} is the fraction of bainite (B) in a structure (%),

{circle around (3)} is the ratio occupied by martensite (M) having anaverage particle diameter of 5 μm or less in ferrite grain boundaries(including grain boundary triple points) (%),

{circle around (4)} is the ratio (a/b) of a C concentration (a) in theferrite phase to a C concentration (b) in the martensite phase, at ¼tbased on the thickness (t) of a base steel sheet,

{circle around (5)} is the ratio (c/d) of an Mn concentration (c) in theferrite phase to an Mn concentration (d) in the martensite phase, at ¼tbased on the thickness (t) of a base steel sheet,

{circle around (6)} is the ratio occupied by ferrite grains having anaverage grain size of 7 μm or greater in the ferrite phase on the basisof the entire area, and

{circle around (7)} is a yield ratio (YS/TS).

As shown in Tables 1 and 2, it can be seen that Inventive Examples 1-7,which satisfy both steel compositions and manufacturing conditionssuggested in the present disclosure, exhibit tensile strengths of450-650 MPa and yield ratios of 0.57 or less, and hole expansion ratiosof 80% or higher are obtained within a tensile strength range of thepresent disclosure.

FIG. 1 shows C and Mn concentration ratios within a depth of 10 μm fromthe surface layer of base steel and at ¼t of the total thickness, whichwere analyzed with a Count Point Sec (CPS) method by Line analysis usingTEM. It can be seen that C and Mn concentration ratios are significantlydecreased at the surface layer compared to a ¼t point.

FIG. 2 shows microstructural observation of a surface layer consideredwithin a depth of 10 μm including the surface layer of base steel and acentral portion using a SEM, and it can be seen that ferrite in thesurface layer is coarse, and particularly, inter-granular martensitephases are significantly reduced.

Conversely, in the case where steel compositions fall within a range ofthe present disclosure but manufacturing conditions are out of range ofthe present disclosure (Comparative Examples 1-5), or where steelcompositions are out of range of the present disclosure (ComparativeExamples 6-8), C and Mn concentration ratios in the surface layer ofbase steel were higher than those at ¼t, or C and Mn concentrationratios in the martensite phase in the surface layer were higher thanthose at ¼t. Therefore, softening in the surface layer of a base steelsheet is hardly expected, and desired mechanical properties in thepresent disclosure could not be obtained. In the case of these steels,it is expected to result in greater potential defects such as ruptureduring formation.

Specifically, in the case of Comparative Example 1 in Inventive Steel 4,the annealing temperature is low, and the fraction of austenite is thuslowered in a dual-phase temperature region. Therefore, in the finalstructure, the fraction of martensite was lowered, and the yield ratiowas thus increased, so that there was a problem in that elongationfinally decreased.

In the case of Comparative Example 2 in inventive steel 5, the annealingtemperature was too high, and the ratio (%) occupied by the martensite(M) phase having an average particle diameter of 5 μm or less in ferritegrain boundaries (including triple points) was thus lowered, resultingin a decrease in elongation. Thus, Comparative Example 2 in InventiveSteel 5 was inappropriate for the present disclosure.

Comparative Steels 1 and 2 fail to satisfy the composition range of thepresent disclosure despite Cr addition for improving hardenability.Therefore, the ratio (%) occupied by the martensite phase in ferritegrain boundaries (including triple points) was lowered, and elongationwas eventually inadequate compared to the target.

As set forth above, according to exemplary embodiments of the presentdisclosure, it is possible to provide a hot-dip galvanized steel sheetand a hot-dip galvannealed steel sheet, which are capable of obtainingexcellent strength and ductility properties, and which are suitable forautomotive body panels requiring high workability.

While exemplary embodiments have been shown and described above, it willbe apparent to those skilled in the art that modifications andvariations could be made without departing from the scope of the presentinvention as defined by the appended claims.

What is claimed is:
 1. A method for manufacturing a hot-dip galvanizedsteel sheet having excellent elongation properties in which a hot-dipgalvanized layer is formed on a surface of a base steel sheet, themethod comprising: preparing a steel slab and thereafter reheating thesteel slab, the steel slab comprising, 0.02-0.08% of carbon (C),1.3-2.1% of manganese (Mn), 0.3% or less of silicon (Si) (excluding 0%),1.0% or less of chromium (Cr) (excluding 0%), 0.1% or less ofphosphorous (P) (excluding 0%), 0.01% or less of sulfur (S) (excluding0%), 0.01% or less of nitrogen (N) (excluding 0%), 0.02-0.06% ofaluminum (sol, Al), 0.2% or less of molybdenum (Mo) (excluding 0%),0.003% or less of boron (B) (excluding 0%), and a remainder of Fe andother inevitable impurities, in wt %; performing finish hot-rolling onthe reheated steel slab in a temperature range of Ar3+50° C.-950° C.,and thereafter coiling the finish hot-rolled steel sheet at 450-700° C.;performing cold-rolling on the coiled hot-rolled steel sheet with areduction ratio of 40-80%, and thereafter performing continuousannealing on the cold-rolled steel sheet in a temperature range of760-850° C.; performing a first cooling on the continuous annealed steelsheet to a temperature range of 630-670° C. at an average cooling rateof 2-8° C./s, and thereafter performing a second cooling on the firstcooled steel sheet to a temperature range of Ms+20° C. to Ms+50° C. atan average cooling rate of 3-10° C./s, wherein Ms is martensite starttemperature; and performing hot-dip galvanizing on the second cooledsteel sheet in a temperature range of 440-480° C., and thereaftercooling the hot-dip galvanized steel sheet to a temperature of Ms−100°C. or lower at an average cooling rate of 4° C./s or higher.
 2. Themethod of claim 1, wherein the base steel sheet of the hot-dipgalvanized steel sheet, has a microstructure including 90% or more offerrite, and a remainder of martensite and bainite, wherein B (%), asdefined by Equation 1, is 3% or less; contains a percentage ofmartensite (M %) having an average particle diameter of 5 μm or lessoccupying space in ferrite grain boundaries (including grain boundarytriple points) of 90% or higher, as defined by Equation 2; and at ¼tbased on the thickness (t) of the sheet, the ratio of a C concentration(a) in the ferrite phase to a C concentration (b) in the martensitephase, i.e., (a)/(b), is 0.7 or less, and the ratio of an Mnconcentration (c) in the ferrite phase to an Mn concentration (d) in themartensite phase, i.e., (c)/(d), is 0.8 or less,B (%)={BA/(MA+BA)}×  100 [Equation 1] where BA is the area occupied bybainite, and MA is the area occupied by martensite,M (%)={M _(gb)/(M _(gb) +M _(in))}×100  [Equation 2] where M_(gb) is theamount of martensite in ferrite grain boundaries, and M_(in) is theamount of martensite within ferrite grains, the martensite having anaverage particle diameter of 5 μm or less.
 3. The method of claim 1,wherein, in the base steel sheet of the hot-dip galvanized steel sheet,the ferrite phase has an average grain size of 4 μm or greater, and thearea occupied by the ferrite phase having an average grain size of 7 μmor greater in the entire ferrite phase is 10% or higher.